Microstructural Evolution and Mechanical Performance of Two Joints of Medium-Mn Stainless Steel with Low- and High-Alloyed Steels

In this work, 2 mm thick medium-Mn austenitic stainless steel (MMn–SS) plates were joined with austenitic NiCr stainless steel (NiCr–SS) and low-carbon steel (LCS) using the gas tungsten arc welding technique. A precise adjustment of the welding process parameters was conducted to achieve high-quality dissimilar joints of MMn–SS with NiCr–SS and LCS. The microstructural evolution was studied using laser scanning confocal and electron microscopes. Secondary electron imaging and electron backscatter diffraction (EBSD) techniques were intensively employed to analyze the fine features of the weld structures. The mechanical properties of the joints were evaluated by uniaxial tensile tests and micro-indentation hardness (HIT). The microstructure of the fusion zone (FZ) in the MMn–SS joints exhibited an austenitic matrix with a small fraction of δ-ferrite, ~6%. The tensile strength (TS) of the MMn–SS/NiCr–SS joint is significantly higher than that of the MMn–SS/LCS joint. For instance, the TSs of MMn–SS joints with NiCr–SS and LCS are 610 and 340 MPa, respectively. The tensile properties of MMn–SS/LCS joints are similar to those of BM LCS, since the deformation behavior and shape of the tensile flow curve for that joint are comparable with the flow curve of LCS. The HIT measurements show that the MMn–SS/NiCr–SS joint is significantly stronger than the MMn–SS/LCS joint since the HIT values are 2.18 and 1.85 GPa, respectively.


Introduction
Medium-Mn stainless steel (MMn-SS) is a promising material for various engineering structural applications, promoting ultra-high strength and moderate ductility [1,2]. Metallurgically, these types of steels, represented in the series 200 stainless steel category, have a metastable austenitic structure, i.e., low stacking fault energy [3][4][5]. Economically, due to the increasing price of nickel, MMn-SS has fascinated researchers in the last decade and is being utilized as a cheaper alternative in the replacement of austenitic NiCr stainless steels (NiCr-SS), especially when serving in moderately corrosive environments [6][7][8].
From the available literature, MMn-SS bearing N shows better mechanical properties than that made of NiCr-SS [3,9]. For instance, Hamada et al. [5] studied the correlation between the mechanical properties of MMn-SS, Type 201, and the grain size evolution markets. Moreover, the mechanical properties of the MMn-SS are significantly better than those of the NiCr-SS, as shown in the above survey. This encourages steel manufacturers to replace NiCr-SS with MMn-SS. Therefore, the dissimilar welding of MMn-SS with the most commonly used steels (NiCr-SS and LCS) is necessary for promoting cost-effective, lighter, and more efficient structural joints.
It is well known that in a modern car body, dissimilar joints of different steel classes with high-strength steel and low-strength steel, e.g., low-alloyed (C-steel) and high-alloyed steels, are processed by welding. Crashworthiness necessitates high-strength and highductility materials in crash-relevant areas. Hence, multi-material designs are achieved in the automotive industry by combining different steels, such as austenitic steel and carbon steel. MMn-SS is becoming increasingly attractive in the automotive industry sector compared to austenitic stainless steel grade AISI 304 due to its superior strength-ductility balance and cost-effectiveness. Hence, the dissimilar welding joining of MMn-SS with conventional austenitic stainless steel (304) and LCS is gaining attention. However, the high Mn content in steel and fusion-welded joints represents a challenge, since it leads to unstable mechanical behavior in the joint. Within this context, this study makes a further contribution to the understanding of the microstructural evolution and mechanical performance of the joints of MMn-SS with NiCr-SS and LCS.
Hence, the present study aims to address the microstructural evolution and mechanical strength of various MMn-SS joints with NiCr-SS and LCS, so as to show the feasibility of fabricating high-strength joints of MMn-SS at a low cost. The microstructural analysis of the weldments was followed by comprehensive observations using different techniques: SE imaging and EBSD. The tensile properties and hardness measurements of the various MMn-SS weldments have been explored to reveal their mechanical strength.

Base Metals (BMs)
The experimental BMs were commercial sheets of steel: MMn-SS, NiCr-SS, and LCS. The steels were received in the form of sheets with a thickness of 2 mm. The sheets were cut into rectangles with dimensions of 100 × 300 mm 2 . The chemical compositions of the BMs were measured by optical emission spectroscopy (model: FOUNDRY-MASTER Pro, OXFORD INSTRUMENTS, High Wycombe, UK). Table 1 illustrates the chemical compositions of the BMs and the filler metals that were utilized during the processing of the various weldments.

Welding Experiment
The metal joining process was conducted by employing the gas tungsten arc welding (GTAW) technique via the ESAB Tig 4300i AC/DC. During welding, the BMs were fixed by tack welds at their corners. Table 2 illustrates welding parameters such as the applied current and voltage. A pure argon shielding gas was used (purity 99.8%) with a flow Materials 2023, 16, 1624 4 of 23 rate of 15 L/min to protect the weldments from oxidation. The heat input was calculated according to Equation (1) [21]: where η represents the arc efficiency of the GTAW process, taken as 0.6 according to BS EN 1011-1 [29]. The symbols (I), (E), and (S) represent the welding current in amperes (A), arc voltage in volt (V), and welding speed in mm/s. Table 2 lists the calculated heat input of the GTAW technique according to the applied welding parameters. The quality of the weld plate was checked via a non-destructive test (radiographic inspection) and reported in our previous article; see Ref. [21]. In general, the quality of the weld plate was relatively good and without solidification cracks. Two different filler metals, ER308L and ER309MoL, were used to form the dissimilar weldments of MMn-SS/NiCr-SS and MMn-SS/LCS, respectively. The tensile strength of both fillers was 610 MPa. It is worth mentioning that the filler ER309MoL bearing a high content of Mo, 2.5 wt.%, was used in the welding joints of LCS to induce a certain fraction of ferrite (3-10%) in the microstructure of the weldment.

Microstructure Characterization
The microstructures of the FZ and heat-affected zone (HAZ) were studied using secondary electron imaging (SE) and electron backscattering diffraction (EBSD) techniques, using field-emission gun scanning electron microscopy (FEG-SEM) (model: Carl Zeiss Ultra plus: Oberkochen, Germany). The EBSD maps were obtained at an accelerating voltage of 15 kV and a working distance of 12-15 mm. The butt joint weldments were sectioned perpendicularly to the welding direction. Typical metallography preparation steps were carried out, starting with mechanical grinding using SiC papers; then, the ground surfaces were polished using a diamond paste suspension and, finally, the specimens were chemically polished using a colloidal suspension of silica. The Thermo-Calc software program (Solna, Sweden) version 2022 with the TCFE9 database was employed to predict the equilibrium phases in the FZ after the welding process.

Mechanical Properties
A micro-indentation tester (CSM instruments: Needham, MA, USA) equipped with a diamond Berkovich indenter was used to measure the indentation hardness (H IT ) of the different zones in the weldment, e.g., FZ, HAZ, and BM. The hardness tests were conducted under an accelerated load from zero to the maximum force (2 N) with a holding time of 10 s; in this stage, the loading behavior was recorded. Subsequently, the applied load was removed with the unloading rate of 66.6 mN/s to record the unloading behavior. The Oliver-Pharr method was automatically applied to analyze the slopes of the loading-unloading curves at the maximum force (2 N) to obtain the H IT values of the weld zones [30].
The mechanical tensile properties of the BMs and the welded joints were evaluated using a hydraulic universal testing machine (UH-F1000kNI SHIMADZU: Kyoto, Japan) at a quasi-static strain rate of 10 −3 s −1 at room temperature. Dog bone-shaped specimens were prepared with a gauge length of 50 mm according to the ASTM standard E8 [31]. Three tensile samples for each dissimilar joint were evaluated. The average tensile results have been provided. Figure 1 shows the microstructures of the experimental materials: BMs of MMn-SS, NiCr-SS, and LCS. Both MMn-SS and NiCr-SS exhibited similar image quality (IQ) maps and phase maps with a blue color, i.e., a fully austenitic γ-fcc structure. The MMn-SS exhibited a fine grain structure with an average grain size of 8.5 ± 4 µm, as shown in Figure 1a. About~5% of low-angle grain boundaries (LAGBs < 15 • ) were promoted in the structure. The corresponding phase map shown in Figure 1b displays a small fraction of α -martensite,~3%, that coexists with the austenitic matrix. This martensite formation in the initial structure of MMn-SS is attributable to the mechanical polishing, which promoted strain, and consequently martensite was induced, as reported in the literature [32,33].  The NiCr-SS exhibited a fully austenitic structure with a grain size of 6 ± 2 µm, which is relatively finer than that of the MMn-SS. A small fraction of LAGBs (~5%) appeared in the grain structure, as displayed in Figure 1c. The corresponding phase map is shown in Figure 1d. The γ-fcc structure was only indexed by EBSD in a blue color.

Microstructure of the BMs
The LCS showed a different phase structure compared with MMn-SS and NiCr-SS, as shown in Figure 1e,f. A ferritic matrix with a grain size of 10 µm is displayed. The LCS contained a small fraction of pearlite, about 7%. However, the pearlite phase was not indexed by EBSD due to the low sensitivity of the EBSD towards the pearlite structure [34]. Hence, the phase map of the LCS exhibits one phase (ferrite), which is represented in red in Figure 1f. A higher fraction of LAGBs,~8%, was present in the grain structure of the LCS, as shown in Figure 1e Figure 2 shows a general view of the microstructure of the FZ in the dissimilar joint of MMn-SS and NiCr-SS at a low magnification, taken via EBSD. The FZ comprised coarse columnar grains and fine equiaxed grains, owing to the heterogeneous thermal distribution during the welding process. The FZ structure mainly comprised coarse columnar grains of size 500 µm near to the BM, and equiaxed grains with an average grain size of~40 µm in the middle of the FZ, ahead of the columnar front, as shown in Figure 2a. This morphology of the weld structure can be attributed to the directional solidification from the edges of the weld pool towards the center [35]. A striking feature of the grain boundaries in the weld structure is the high presence of LAGBs,~25%, as indicated by the green boundaries in Figure 2a. size of ∼40 μm in the middle of the FZ, ahead of the columnar front, as shown in Figure  2a. This morphology of the weld structure can be attributed to the directional solidification from the edges of the weld pool towards the center [35]. A striking feature of the grain boundaries in the weld structure is the high presence of LAGBs, ∼25%, as indicated by the green boundaries in Figure 2a.   Figure 2c shows that the phase structure of the FZ was mainly austenite (γ-fcc) with a small fraction of ferrite (δ-bcc), while the volume fraction of the δ-ferrite was approximately ∼5.7%, as measured by EBSD. The grain size distribution of the grain structure of the FZ is shown in Figure 2d. It can be   Figure 2c shows that the phase structure of the FZ was mainly austenite (γ-fcc) with a small fraction of ferrite (δ-bcc), while the volume fraction of the δ-ferrite was approximately~5.7%, as measured by EBSD. The grain size distribution of the grain structure of the FZ is shown in Figure 2d. It can be seen that 90% of the grains had sizes smaller than 100 µm. Few grains had coarse sizes of 200-800 µm. Figure 3 shows a magnified view of the FZ microstructure at the middle zone of the grain structure. Figure 3a displays a dual-phase structure: γ-austenite + δ-ferrite. Sabzi et al. [36] reported a similar dual-phase structure of the weld zone in the joint of two alloys of stainless steel (316L/310S). The IQ map shows various grey levels for the austenitic grains owing to the different crystallographic textures of the grains. The regularity of the δ-ferrite distribution is a conspicuous feature of the weld structure, since it was promoted with two morphologies: ferrite laths at austenite grain boundaries and vermicular ferrite within the austenitic grains, as shown in the phase map, Figure 3b. A detailed investigation of the morphologies of the austenite and ferrite, undertaken via SEM, is presented in Section 4. δ-ferrite distribution is a conspicuous feature of the weld structure, since it was promoted with two morphologies: ferrite laths at austenite grain boundaries and vermicular ferrite within the austenitic grains, as shown in the phase map, Figure 3b. A detailed investigation of the morphologies of the austenite and ferrite, undertaken via SEM, is presented in Section 4. Figure 4 mainly shows the grain structures of three zones related to the BM NiCr-SS, e.g., the HAZ, transition zone (TZ), and fusion boundary (FB), which is a partially melted zone (PMZ) at the interface between the FZ and the BM, highlighted by the yellow dashed lines. Homogeneity in the sizes and grains of the FB along the weld can be observed. Another zone in the fine grain structure, the TZ, highlighted by the white dashed line, appeared between the FB and the HAZ. The mechanism promoting TZs in the stainless-steel BMs is discussed in Section 4.
The width of the FB was relatively small compared to the other zones. For instance, the sizes of the HAZ, TZ, and FB were ∼1400, 350, and 225 μm, respectively. No phase transformation occurred since the austenite was the dominant structure in the FB, TZ, and HAZ. It is noteworthy that the annealing twins disappeared in the FB, while they remained noticeable in the TZ and HAZ. The twin plates were reoriented in a similar manner as the parent matrix when the temperature reached 1000 °C (known as the de-twinning phenomenon) [37,38]. Therefore, the disappearance of the twin plates in the FB confirms that the temperature there exceeded 1000 °C, while the existence of the twin plates in the TZ confirms that the temperature there was less than 1000 °C.  4 mainly shows the grain structures of three zones related to the BM NiCr-SS, e.g., the HAZ, transition zone (TZ), and fusion boundary (FB), which is a partially melted zone (PMZ) at the interface between the FZ and the BM, highlighted by the yellow dashed lines. Homogeneity in the sizes and grains of the FB along the weld can be observed. Another zone in the fine grain structure, the TZ, highlighted by the white dashed line, appeared between the FB and the HAZ. The mechanism promoting TZs in the stainless-steel BMs is discussed in Section 4.
The width of the FB was relatively small compared to the other zones. For instance, the sizes of the HAZ, TZ, and FB were~1400, 350, and 225 µm, respectively. No phase transformation occurred since the austenite was the dominant structure in the FB, TZ, and HAZ. It is noteworthy that the annealing twins disappeared in the FB, while they remained noticeable in the TZ and HAZ. The twin plates were reoriented in a similar manner as the parent matrix when the temperature reached 1000 • C (known as the de-twinning phenomenon) [37,38]. Therefore, the disappearance of the twin plates in the FB confirms that the temperature there exceeded 1000 • C, while the existence of the twin plates in the TZ confirms that the temperature there was less than 1000 • C. The grain structure of the FB consisted of columnar grains oriented towards the BM. This is attributed to the unidirectional heat flux from the molten pool in the FZ towards the BM, as reported in the literature [39]. However, fine equiaxed grains can be observed in the TZ. The transition of the columnar structure in the FB into fine equiaxed grains in the TZ is discussed in Section 4. As expected, a coarse grain structure was promoted in the HAZ, as shown in Figure 4a. The corresponding IPF map shown in Figure 4b indicates that the grains in these zones preferred <101> and <111> crystallographic orientations.  Figure 5a shows the IQ map of the grain structure, which displays irregular and columnar grains. Meanwhile, Figure 5a shows that a significant fraction of LAGBs, ∼42%, was induced in the weld structure. This is presumably attributable to the residual stresses promoted in the welding during fast cooling [40,41]. The corresponding IPF map shown in Figure 5b indicates that the orientation of the grains in the FZ preferred <101> and <001> crystallographic orientations. The phase map in Figure 5c shows a dual-phase structure, since the structure is mainly austenite with a few fractions of δ-ferrite, about 3.5%, as calculated by EBSD. The grain structure of the FB consisted of columnar grains oriented towards the BM. This is attributed to the unidirectional heat flux from the molten pool in the FZ towards the BM, as reported in the literature [39]. However, fine equiaxed grains can be observed in the TZ. The transition of the columnar structure in the FB into fine equiaxed grains in the TZ is discussed in Section 4. As expected, a coarse grain structure was promoted in the HAZ, as shown in Figure 4a. The corresponding IPF map shown in Figure 4b indicates that the grains in these zones preferred <101> and <111> crystallographic orientations.  Figure 5a shows the IQ map of the grain structure, which displays irregular and columnar grains. Meanwhile, Figure 5a shows that a significant fraction of LAGBs,~42%, was induced in the weld structure. This is presumably attributable to the residual stresses promoted in the welding during fast cooling [40,41]. The corresponding IPF map shown in Figure 5b indicates that the orientation of the grains in the FZ preferred <101> and <001> crystallographic orientations. The phase map in Figure 5c shows a dual-phase structure, since the structure is mainly austenite with a few fractions of δ-ferrite, about 3.5%, as calculated by EBSD.   Figure 6a. White thin precipitates of the δ-bcc phase with an acicular morphology were mainly located in the inter-dendritic regions of the austenite phase. This is confirmed by the corresponding phase map in Figure 6b, which depicts the matrix austenitic grains in blue and the δ-ferrite in red. It can be observed that δ-ferrite with different morphologies is promoted inside the austenitic grains, as highlighted by the yellow dashed ovals. Similarly, Chuaiphan et al. [42] reported a duplex structure of austenite and δ-ferrite in the microstructures of the FZ in the joint of 304-SS and LCS grade 1020 welded via the GTAW technique.   Figure 6a. White thin precipitates of the δ-bcc phase with an acicular morphology were mainly located in the inter-dendritic regions of the austenite phase. This is confirmed by the corresponding phase map in Figure 6b, which depicts the matrix austenitic grains in blue and the δ-ferrite in red. It can be observed that δ-ferrite with different morphologies is promoted inside the austenitic grains, as highlighted by the yellow dashed ovals. Similarly, Chuaiphan et al. [42] reported a duplex structure of austenite and δ-ferrite in the microstructures of the FZ in the joint of 304-SS and LCS grade 1020 welded via the GTAW technique. Figure 7 shows the microstructures of the FZ, PMZ (FB), TZ, HAZ, and BM on the MMn-SS side via EBSD. Figure 7a shows the IQ map of the grain structure. Irregular and columnar grains are shown. Meanwhile, Figure 7b displays the corresponding IPF map of the orientation of the grains in the FZ and HAZ, exhibiting different crystallographic orientations. Figure 7c displays the corresponding phase map, wherein the structure of the HAZ is shown to have been mainly austenite.
The microstructure of the MMn-SS displays an FB with a width of 0.216 ± 0.06 mm. Moreover, the TZ showed a width of 0.424 ± 0.1 mm, while the HAZ showed a width of 1.8 ± 0.3 mm. The width of the coarse grain HAZ (CGHAZ) was 1.1 ± 0.1 mm, with an average grain size of 60 µm, while the width of the fine grain HAZ (FGHAZ) was 0.7 ± 0.1 mm, with an average grain size of 25 µm. Figure 8 shows the microstructures of FZ, PMZ, HAZ, and BM on the LCS side, taken via EBSD. Figure 8a displays an IQ map of the grain structure. The columnar grain structure was limited at the CGHAZ on the LCS side, while an equiaxed structure was promoted. Figure 8b exhibits the corresponding IPF map of the orientation of the grains in the FZ and HAZ, with different crystallographic orientations. Figure 8c shows the corresponding phase map, wherein it is clear that the structure in the HAZ is mainly α-ferrite (colored in red).
Furthermore, the LCS did not show a TZ such as that found after the welding of the stainless steel joints. Moreover, the microstructure of the LCS showed a very thin PMZ with a width of 0.016 ± 0.008 mm. The HAZ showed a width of 2 ± 0.15 mm. The average grain size in CGHAZ was measured at 55 ± 25 µm, while the average grain size in the FGHAZ was measured at 20 ± 10 µm.

Mechanical Properties of MMn-SS Weld Joints
The engineering tensile stress-strain curves of the MMn-SS joints with NiCr-SS and LCS at room temperature are shown in Figure 9. A striking feature of the tensile flow behavior is the considerable deformation capacity and strain hardening of the MMn-SS/NiCr-SS joint. During tensile testing, the joint was broken within the weld. Hence, the yield strength and UTS of the joint, 340 and 610 MPa, were relatively lower than those of the BMs (MMn-SS and NiCr-SS). However, the tensile flow behavior of the joint MMn-SS/LCS resembles the tensile flow behavior of the LCS: a Lüders plateau with a low yield strength. This is because the joint was broken within the LCS BM.
The mechanical tensile properties of the MMn-SS joints compared with the BMs are illustrated in Table 3. It can be observed that the mechanical strengths of the BMs, MMn-SS, and NiCr-SS were higher than those of the weldment; for instance, the UTSs of the BMs and weld were 915, 738, and 610 MPa, respectively. However, the mechanical strength of the other joint with LCS was slightly lower than that of the LCS BM, 340 and 360 MPa, respectively.   Figure 7a shows the IQ map of the grain structure. Irregular and columnar grains are shown. Meanwhile, Figure 7b displays the corresponding IPF map of the orientation of the grains in the FZ and HAZ, exhibiting different crystallographic orientations. Figure 7c displays the corresponding phase map, wherein the structure of the HAZ is shown to have been mainly austenite.
The microstructure of the MMn-SS displays an FB with a width of 0.216 ± 0.06 mm.   Figure 8a displays an IQ map of the grain structure. The columnar grain structure was limited at the CGHAZ on the LCS side, while an equiaxed structure was promoted. Figure 8b exhibits the corresponding IPF map of the orientation of the grains in the FZ and HAZ, with different crystallographic orientations. Figure 8c shows the corresponding phase map, wherein it is clear that the structure in the HAZ is mainly α-ferrite (colored in red).     Figure 10 shows the fracture surfaces of the welded joints after tensile testing. Figure  10a,b show the fractography of the tensile specimens, MMn-SS/NiCr-SS and MMn-SS/LCS, respectively. It can be observed that the fracture surfaces of the welded joints exhibited a dimple feature with varying shapes and sizes, which is a typical characteristic of a ductile fracture.  It is well known that the weld efficiency is related to the tensile strength of the processed joint. Equation (2) shows the terms for calculating the weld efficiency (WE) in both joints [43]: where TS WM is the UTS of the weld, and TS BM is the UTS of the weaker BM in the joint. The WE parameter of the welded joints was calculated based on the tensile properties extracted from the tensile flow curves (Figure 9). The WEs of the MMn-SS/NiCr-SS and MMn-SS/LCS joints were 82 and 95%, respectively; while the MMn-SS/NiCr-SS joint was slightly softer than the NiCr-SS BM, the MMn-SS/LCS joint exhibited a similar strength to the LCS-BM. Figure 10 shows the fracture surfaces of the welded joints after tensile testing. Figure 10a,b show the fractography of the tensile specimens, MMn-SS/NiCr-SS and MMn-SS/LCS, respectively. It can be observed that the fracture surfaces of the welded joints exhibited a dimple feature with varying shapes and sizes, which is a typical characteristic of a ductile fracture.

Discussion
In the present study, the dissimilar joints of MMn-SS/NiCr-SS and MMn-SS/LCS were processed by the GTAW technique via ER308L and ER309MoL filler metals, at heat inputs of 0.486 and 0.604 kJ/mm, respectively. In this section, we explain variations in the micro-indentation hardness values in view of the microstructural evolution in the FZ and TZ of the welded joints. Figure 11 shows the Isopleth equilibrium phase diagrams for the chemical composition of the FZs, analyzed by EDS at several spots in the studied joints using the TCFE9 Thermo-Calc software. The volume fractions of the equilibrium phases are displayed in the temperature range of 1500 to 500 °C. It can be observed that the δ-ferrite (red line) started to form from the molten pool at 1450 and 1410 °C in the joints of MMn-SS/NiCr-SS ( Figure 11a) and MMn-SS/LCS (Figure 11b), respectively. The proportions of δ-ferrite were ~97% and ~80%, mixed with 3% and 20% liquid, in the FZ of the MMn-SS/NiCr-SS and MMn-SS/LCS joints, respectively. This implies the possibility of δ-ferrite forming at a higher fraction in the MMn-SS/NiCr-SS than in the MMn-SS/LCS joint, which depends on the cooling rate after weld processing (that is related to the heat input consumed during the welding process) [21].

Thermo-Calc Analysis of the Predicted Phases in the FZ
With cooling, the δ-ferrite and the existing liquid transformed into an austenite phase (blue line) up to complete transformation at 1150 and 1120 °C in the joints MMn-SS/NiCr-SS and MMn-SS/LCS, respectively. With cooling to the lower temperature of 500 °C, αferrite formed at fractions of 50% and 10% in the FZs of MMn-SS/NiCr-SS and MMn-SS/LCS, respectively. However, these phases were in equilibrium. It is well established that the welding process involves a non-equilibrium state, due to the fast cooling after the welding processing. Hence, the δ-ferrite formed at high temperatures was kept in the structures of the weldments. The Thermo-Calc plots show the formation of carbide-nitride precipitates due to the existence of the alloying elements, see Section 4.2.

Discussion
In the present study, the dissimilar joints of MMn-SS/NiCr-SS and MMn-SS/LCS were processed by the GTAW technique via ER308L and ER309MoL filler metals, at heat inputs of 0.486 and 0.604 kJ/mm, respectively. In this section, we explain variations in the micro-indentation hardness values in view of the microstructural evolution in the FZ and TZ of the welded joints. Figure 11 shows the Isopleth equilibrium phase diagrams for the chemical composition of the FZs, analyzed by EDS at several spots in the studied joints using the TCFE9 Thermo-Calc software. The volume fractions of the equilibrium phases are displayed in the temperature range of 1500 to 500 • C. It can be observed that the δ-ferrite (red line) started to form from the molten pool at 1450 and 1410 • C in the joints of MMn-SS/NiCr-SS ( Figure 11a) and MMn-SS/LCS (Figure 11b), respectively. The proportions of δ-ferrite werẽ 97% and~80%, mixed with 3% and 20% liquid, in the FZ of the MMn-SS/NiCr-SS and MMn-SS/LCS joints, respectively. This implies the possibility of δ-ferrite forming at a higher fraction in the MMn-SS/NiCr-SS than in the MMn-SS/LCS joint, which depends on the cooling rate after weld processing (that is related to the heat input consumed during the welding process) [21]. It is well established that different solidification modes occur after the weld processing of austenitic steels based on the ratio of Creq/Nieq. For instance, the ferritic austenitic (FA) mode refers to the solidification of ferrite, followed by austenite formation, which is promoted when 1.48 < (Creq/Nieq) < 1.95 [11,23].

Thermo-Calc Analysis of the Predicted Phases in the FZ
In the present study, the Creq and Nieq values were calculated according to Equations With cooling, the δ-ferrite and the existing liquid transformed into an austenite phase (blue line) up to complete transformation at 1150 and 1120 • C in the joints MMn-SS/NiCr-SS and MMn-SS/LCS, respectively. With cooling to the lower temperature of 500 • C, α-ferrite formed at fractions of 50% and 10% in the FZs of MMn-SS/NiCr-SS and MMn-SS/LCS, respectively. However, these phases were in equilibrium. It is well established that the welding process involves a non-equilibrium state, due to the fast cooling after the welding processing. Hence, the δ-ferrite formed at high temperatures was kept in the structures of the weldments. The Thermo-Calc plots show the formation of carbide-nitride precipitates due to the existence of the alloying elements, see Section 4.2.
It is well established that different solidification modes occur after the weld processing of austenitic steels based on the ratio of Cr eq /Ni eq . For instance, the ferritic austenitic (FA) mode refers to the solidification of ferrite, followed by austenite formation, which is promoted when 1.48 < (Cr eq /Ni eq ) < 1.95 [11,23].
In the present study, the Cr eq and Ni eq values were calculated according to Equations (3) and (4), as follows: Cr eq = Cr + Mo + 0.5 Nb + 1.5 Si Ni eq = Ni + 30 (C + N) + 0.5 (Mn + Cu) The chemical compositions of the FZs were determined by EDS attached to an SEM to calculate Cr eq and Ni eq . On the one hand, the EDS has a low sensitivity in the detection of low-atomic number elements, such as C and N, since the sensitivity of EDS analysis is approximately 0.1 wt.% for all elements [44,45]. On the other hand, the calculation of the elements present at the FZ by EDS is accepted in numerous other works [11,13,17]. The C content in the FZs was calculated by the stoichiometry method, which depends on the quantitative relationship between the C contents of the paired steel plates (MMn-SS/NiCr-SS and MMn-SS/LCS) and the filler metals (ER308L and ER309MoL), as well as the effect of the dilution percentage for each steel in the joint. This is determined according to Equation (5), as follows [11]: C FZ = (C.D) BM1 + (C.D) BM2 + (C.D) FM (5) where C FZ is the C content of the FZ, (C.D) BM1 represents the C content and dilution percentage of BM 1, (C.D) BM2 is the C content and dilution percent of BM 2, and (C.D) FM represents the C content and dilution percent of the filler metal utilized. Based on the δ-ferrite content measured via EBSD, dilution percentages of 15% and 70% were assumed for the paired steels and the fillers, respectively. Consequently, the C contents in the FZs were 0.032 and 0.029 wt.% for MMn-SS/NiCr-SS and MMn-SS/LCS, respectively.
Here, the ratio of Cr eq /Ni eq was 1.8 and 1.53 for MMn-SS/NiCr-SS and MMn-SS/LCS, respectively. This indicates that the weld pool exhibited an FA solidification mode, which means that the δ-ferrite was solidified first, as it preceded the formation of the austenite phase, and this agrees with the results of the Thermo-Calc shown in Figure 11.
The δ-ferrite volume fraction existing in the FZ of the MMn-SS/LCS joint (3.5%) was lower than that existing in the FZ of the MMn-SS/Ni-Cr-SS joint, which was 5.7%. Despite this, the filler, ER308L, utilized during the welding of the MMn-SS/NiCr-SS, had a Mo content (0.2 wt.%) lower than that of the filler ER309MoL (2.5 wt.%) used in the welding of the MMn-SS/LCS joint. It is well known that Mo is a ferrite stabilizer, as reported in the literature [46,47].
It is reasonable to assume that the change in the fractions of δ-ferrite present in the FZs of the MMn-SS/NiCr-SS and MMn-SS/LCS joints can be attributed to two factors. First, the heat input in the MMn-SS/NiCr-SS joint was lower than that in the MMn-SS/LCS joint, 0.486 compared to 0.604 kJ/mm, respectively. It is established that the higher the heat input, the lower the cooling rate [21]. Therefore, the δ-ferrite (which was solidified first, according to the FA solidification mode) was isothermally transformed into austenite in the FZ of the MMn-SS/LCS joint due to the relatively slow cooling rate compared to the cooling rate of the MMn-SS/NiCr-SS. Abioye et al. [48] reported that the volume fraction of the δ-ferrite phase increases as the heat input decreases during the dissimilar welding of NiCr-SS (304) with LCS. Second, the high alloying element, Cr, in the FZ of the MMn-SS/NiCr-SS, promotes a higher δ-ferrite content than that in the MMn-SS/LCS joint. Figure 12 shows the dendritic structure of the FZ, determined by laser microscopy and SEM imaging. A conspicuous feature of the grain structure in the FZ can be observed, i.e., the irregularity of the morphology of the grains, as shown in Figure 12a. For instance, the columnar austenitic dendritic morphology, highlighted in red, and the equiaxed cellular grains, highlighted in the yellow oval, are promoted in the FZ. It can be confirmed that the FZ structure was composed of γ-austenite and δ-ferrite phases, as seen in magnified views taken by SEM; Figure 12b-d. Figure 12c reveals a "Widmanstätten pattern" in the present phases, i.e., lamellar austenite grains and parallel acicular ferrite located at the austenite grain boundaries. Ferrite was present with two morphologies, ferrite laths at the austenite grain boundaries and vermicular ferrite within the austenitic grains, similar to what has been reported in the literature [6]. Figure 12d shows a higher magnification of the FZ structure, exhibiting that other parts of the δ-ferrite were located inside the bulk of the austenite grains and had branched formations. The Widmanstätten ferrite lamellas were very thin, <1 µm (which agrees with the results presented by EBSD in Figure 3b).

Characteristics of δ-Ferrite in the FZs
It is reasonable to assume that the change in the fractions of δ-ferrite present in the FZs of the MMn-SS/NiCr-SS and MMn-SS/LCS joints can be attributed to two factors. First, the heat input in the MMn-SS/NiCr-SS joint was lower than that in the MMn-SS/LCS joint, 0.486 compared to 0.604 kJ/mm, respectively. It is established that the higher the heat input, the lower the cooling rate [21]. Therefore, the δ-ferrite (which was solidified first, according to the FA solidification mode) was isothermally transformed into austenite in the FZ of the MMn-SS/LCS joint due to the relatively slow cooling rate compared to the cooling rate of the MMn-SS/NiCr-SS. Abioye et al. [48] reported that the volume fraction of the δ-ferrite phase increases as the heat input decreases during the dissimilar welding of NiCr-SS (304) with LCS. Second, the high alloying element, Cr, in the FZ of the MMn-SS/NiCr-SS, promotes a higher δ-ferrite content than that in the MMn-SS/LCS joint. Figure 12 shows the dendritic structure of the FZ, determined by laser microscopy and SEM imaging. A conspicuous feature of the grain structure in the FZ can be observed, i.e., the irregularity of the morphology of the grains, as shown in Figure 12a. For instance, the columnar austenitic dendritic morphology, highlighted in red, and the equiaxed cellular grains, highlighted in the yellow oval, are promoted in the FZ. It can be confirmed that the FZ structure was composed of γ-austenite and δ-ferrite phases, as seen in magnified views taken by SEM; Figure 12b-d. Figure 12c reveals a "Widmanstätten pattern" in the present phases, i.e., lamellar austenite grains and parallel acicular ferrite located at the austenite grain boundaries. Ferrite was present with two morphologies, ferrite laths at the austenite grain boundaries and vermicular ferrite within the austenitic grains, similar to what has been reported in the literature [6]. Figure 12d shows a higher magnification of the FZ structure, exhibiting that other parts of the δ-ferrite were located inside the bulk of the austenite grains and had branched formations. The Widmanstätten ferrite lamellas were very thin, <1 μm (which agrees with the results presented by EBSD in Figure 3b).   Figure 13a,b display intermittent δ-ferrite at the grain boundaries of the austenite, as well as precipitates of carbides, which appeared randomly at the grain boundaries and inside the grains of the austenite phase. Figure 13c exhibits the elemental distribution map of the elements present in the microstructure; the map shows darker colors in intermittent forms on the grain boundaries at the locations of the δ-ferrite. The presence of the δ-ferrite phase with an intermittent shape confirms that the transformation of the δ-ferrite into austenite was enhanced due to the slow cooling rate associated with the relatively high heat input applied during the welding processing. This agrees with the findings in the literature [11], which reported that with the rise in the cooling rate, lower percentages of δ-ferrite are transformed into austenite. This result highlights the serious effects of the heat input entered during the welding of MMn-SS joints, since it directly affects the amount of δ-ferrite present in the FZ. Figure 13 shows magnified views of the FZ in the MMn-SS/LCS welded joint, a map analysis of the present elements, and the corresponding EDS analysis performed via SEM. Figure 13a,b display intermittent δ-ferrite at the grain boundaries of the austenite, as well as precipitates of carbides, which appeared randomly at the grain boundaries and inside the grains of the austenite phase. Figure 13c exhibits the elemental distribution map of the elements present in the microstructure; the map shows darker colors in intermittent forms on the grain boundaries at the locations of the δ-ferrite. The presence of the δ-ferrite phase with an intermittent shape confirms that the transformation of the δ-ferrite into austenite was enhanced due to the slow cooling rate associated with the relatively high heat input applied during the welding processing. This agrees with the findings in the literature [11], which reported that with the rise in the cooling rate, lower percentages of δ-ferrite are transformed into austenite. This result highlights the serious effects of the heat input entered during the welding of MMn-SS joints, since it directly affects the amount of δ-ferrite present in the FZ.

Characteristics of δ-Ferrite in the FZs
The EDS analysis (Figure 13d) shows that the δ-ferrite contained high amounts of Cr (~20 wt.%). On the other hand, the precipitated carbides on the grain boundaries contained C and Cr, while the carbides precipitated inside the austenite matrix contained C, Cr, and Mo. The EDS analysis of spot #3, shown in Figure 13d, displays that the microscale pitting/porosity on the surface was promoted in the FZ, which was related to the high heat input employed during the welding, and this is consistent with the literature [49]. The EDS analysis (Figure 13d) shows that the δ-ferrite contained high amounts of Cr (~20 wt.%). On the other hand, the precipitated carbides on the grain boundaries contained C and Cr, while the carbides precipitated inside the austenite matrix contained C, Cr, and Mo. The EDS analysis of spot #3, shown in Figure 13d, displays that the micro-scale pitting/porosity on the surface was promoted in the FZ, which was related to the high heat input employed during the welding, and this is consistent with the literature [49]. Figures 4 and 7 show the appearance of the TZ between the FBs and HAZs of the NiCr-SS and MMn-SS sides, while Figure 8 does not indicate the existence of any TZ on the LCS side. As reported in the Results section, the grain size in the TZ was finer than the grain size of the neighboring CGHAZ. Therefore, a detailed microstructural investigation of the TZ in the MMn-SS side is presented in Figure 14, which explains the formation of the fine grain structure there.

Evolution of Micro-Indentation Hardness Values
Micro-indentation hardness tests of the welded joints were carried out to evaluate the effect of the present phases in the microstructure, as well as grain size variations, on the evolution of the hardness in the welded joints. Figure 15 presents the welded joints' typical micro-indentation load-penetration depth (P-h) curves for different zones (BM, HAZ, and FZ). The penetration load was kept constant at 2 N. The variation in penetration depths across the same weldments is attributed to the microstructural and compositional inhomogeneity.
The BM MMn-SS recorded a minimum penetration depth of 5.4 ± 0.2 μm (corresponding to HIT = 3.16 ± 0.13 GPa), while the BM NiCr-SS recorded a penetration depth of 5.75 ± 0.1 μm (corresponding to HIT = 2.81 ± 0.01 GPa). Despite the NiCr-SS showing a grain size (6 μm) that was finer than that of the MMn-SS (8.5 um), the MMn-SS had a higher content of interstitial atoms, such as C and N. For instance, the C and N contents in the MMn-SS BM were 0.1 and 0.2 wt.%, respectively, while the corresponding C and N contents in the BM NiCr-SS were 0.03 and 0.1 wt.%, respectively. Furthermore, the BM MMn-SS had a higher tensile strength, 915 MPa, than that of the NiCr-SS, 738 MPa. Therefore, the penetration depth in the MMn-SS was smaller than that in the NiCr-SS. On the other hand, the LCS-BM recorded a penetration depth of 7.15 ± 0.1 μm (corresponding to HIT = 1.6 ± 0.05 GPa). These results confirm that the ferrite phase had the lowest resistance to indentation among the present phases.
The penetration depth in the FZ of the stainless steel weldment, MMn-SS/NiCr-SS,  Figure 14a shows a formation of longitudinal laths or fine precipitates along the grain boundaries of the austenite matrix in the TZ. Figure 14b confirms that those precipitates did not appear at the CGHAZ, while a grain coarsening effect was promoted. Figure 14c displays a magnified view of the TZ; it shows that the precipitates are located at the grain boundaries of the austenite phase. Furthermore, twin plates appear inside the austenite grains; they were stable despite the heat passing from the FZ towards the BM, which confirms that the temperature in the TZ was less than 1000 • C [37,38].
The main reason behind the appearance of the TZ is the sensitization effect [50][51][52]. The sensitization takes place in the HAZ region at a temperature range of 400 to 800 • C after the welding of austenitic steels. It refers to the Cr-depletion of the zones close to the grain boundaries, which is followed by a formation of Cr-rich chains at the grain boundaries of the austenite matrix, whereas C is diffused out of the austenite grains to enhance the formation of C-Cr precipitates along the grain boundaries. However, the formation of the C-Cr-rich chains along the grain boundaries negatively affects the corrosion behavior of austenitic steels [53]. Furthermore, according to the results presented in Figure 11, chromium-carbides started to precipitate at 850 and 950 • C for MMn-SS/NiCr-SS and MMn-SS/LCS, respectively. Figure 14d exhibits the chemical analysis of the precipitates (located at the grain boundaries) via the EDS; it confirms that they are Chromium carbides, as it mainly contains C, Cr, and Mn elements. It was reported that the precipitation of carbide particles on the grain boundaries has pinning effects on the grain growth [54][55][56]. Therefore, the precipitated carbides in the TZ work as obstacles for the movement of the grain boundaries, which hinders the coarsening of the grains size at the TZ and works as heterogeneous nucleation sites, which enhances the formation of fine equiaxed structures in the TZ; therefore, the size of the grains at the TZ was finer than the grains which exist at the FB and CGHAZ, which is similar to the results presented in the literature [52,57,58].

Evolution of Micro-Indentation Hardness Values
Micro-indentation hardness tests of the welded joints were carried out to evaluate the effect of the present phases in the microstructure, as well as grain size variations, on the evolution of the hardness in the welded joints. Figure 15 presents the welded joints' typical micro-indentation load-penetration depth (P-h) curves for different zones (BM, HAZ, and FZ). The penetration load was kept constant at 2 N. The variation in penetration depths across the same weldments is attributed to the microstructural and compositional inhomogeneity. SS/LCS joint showed a penetration depth of 6.7 ± 0.15 μm (corresponding to HIT = 1.85 ± 0.1 GPa), which was softer than the hardness of the FZ of the MMn-SS/NiCr-SS joint (2.18 ± 0.1 GPa). However, the FZ in the Mn-SS/NiCr-SS joint exhibited a higher fraction of the soft phase of δ-ferrite, which reduced the hardness, as reported in Ref. [59]. However, the higher HIT of the FZ in the Mn-SS/NiCr-SS joint was mainly attributed to the solid solution strengthening mechanism, related to the high presence of alloying elements (Ni, Cr, Fe) promoted by combing three high-alloyed steels (MMn-SS, NiCr-SS, filler ER308L). This is in agreement with the findings of Karimi et al. [60]. Furthermore, the HAZ of the LCS side exhibited the highest penetration depth of 7.55 ± 0.1 μm, corresponding to HIT = 1.45 ± 0.05 GPa. This can be attributed to the coarsening of the grain structure in HAZ.

Conclusions
The joints of medium-Mn stainless steel with a high tensile strength of 915 MPa, NiCr stainless steel with a strength of 738 MPa, and low C-steel with a low tensile strength of 360 MPa were manufactured using the GTAW technique to assess the durability and feasibility of using low-cost joints for structural applications. Microstructural and mechanical performance (tensile strength and micro-indentation hardness) analyses of the dissimilar metal joints were performed. The main conclusions can be summarized based on the intensive experimental observations as follows: